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www.elsevier.com/locate/surfcoat
Surface & Coatings Technolog
Abrasive wear behaviour of WC–CoCr and Cr3C2–20(NiCr) deposited by
HVOF and detonation spray processes
J.K.N. Murthy, B. Venkataraman*
Defence Metallurgical Research Laboratory, P.O. Kanchanbagh, Hyderabad-500048, India
Received 18 June 2004; accepted in revised form 28 October 2004
Available online 24 December 2004
Abstract
Thermally sprayed tungsten carbide-based and chromium carbide-based coatings are being widely used for a variety of wear resistance
applications. These coatings deposited by high velocity processes like high velocity oxy-fuel (HVOF) and detonation gun spray (DS)
techniques are known to provide improved wear performance. In the present study, WC–10Co–4Cr and Cr3C2–20(NiCr) coatings are
deposited by HVOF and pulsed DS processes, and low stress abrasion wear resistance of these coatings are compared. The abrasion tests
were done using a three-body solid particle rubber wheel test rig using silica grits as the abrasive medium. The results show that the DS
coating performs slightly better than the HVOF coating possibly due to the higher residual compressive stresses induced by the former
process and WC-based coating has higher wear resistance in comparison to Cr3C2-based coating. Also, the thermally sprayed carbide-based
coatings have excellent wear resistance with respect to the hard chrome coatings.
D 2004 Elsevier B.V. All rights reserved.
Keywords: WC–CoCr; Cr3C2–20(NiCr); HVOF and detonation spray processes
1. Introduction
Thermally sprayed cermet coatings have emerged as a
viable solution for a wide range of wear resistance
applications to improve the service life of machine
components. Tungsten carbide and chromium carbide-based
coatings are frequently used for many of the applications in
gas turbine, steam turbine and aero-engine to improve the
resistance to sliding, abrasive and erosive wear [1,2]. The
former is used up to 500 8C and the latter up to 800 8C[3,4]. Also, for sliding wear and abrasive wear resistance,
the carbide coatings are considered to be a viable
alternative to hard chrome platings due to the strict
environmental regulations and cost concerns with regard
to the electroplating process [5,6]. These cermet coatings
0257-8972/$ - see front matter D 2004 Elsevier B.V. All rights reserved.
doi:10.1016/j.surfcoat.2004.10.136
* Corresponding author. Tel.: +91 40 24586476; fax: +91 40 24340683/
24341439.
E-mail addresses: [email protected],
[email protected] (B. Venkataraman).
are deposited by plasma spray and high velocity processes
namely high velocity oxy-fuel (HVOF) and detonation gun
spray (DS) processes. The high velocity processes namely
the HVOF and DS are usually employed for depositing
these coatings to avoid significant amount of reduction of
carbides to brittle carbides and oxy-carbides due to the
much lower temperature of the powder particles in the
exhaust gas stream and less in-flight time as compared to
that in plasma [7,8]. Also, the higher particle velocities in
the high velocity processes lead to better coating properties
like higher bond strength, density and lower oxide content.
It has been reported that carbide containing coatings
deposited by high velocity processes have good wear
resistance [9] compared to plasma-sprayed coatings due to
the better coating properties achievable in case of high
velocity processes as mentioned earlier. WC and Cr3C2
with different metallic binders like Co, Ni and Fe have been
studied using different amounts of binder contents with Co
and Ni most commonly used. Addition of Cr to the matrix
has been found to improve the wear and oxidation
resistance of these cermets [8,10]. The wear behaviour of
y 200 (2006) 2642–2652
Table 2
Spraying conditions adopted for HVOF and DS processes
Coating process Coating
SM 5847 D-3007
1. HVOF process
Spray gun DJ2600 DJ2600
O2 pressure (MPa) 1.14, 1.8 1.17, 1.92
O2 Flow rate (m3/h)
H2 pressure (MPa) 0.93, 3.6 0.96, 3.84
H2 Flow rate (m3/h)
Air pressure (MPa) 0.69, 1.2 0.55, 1.68
Air Flow rate (m3/h)
Spray rate (kg/h) 3.6 3.6
Carrier gas Nitrogen Nitrogen
Spraying distance (m) 0.2 0.2
2. DS process
O2/acetylene volume ratio 1:1.23 1:1.21
Carrier gas flow rate (m3/h) 1.6 1.2
Spray distance (m) 0.165 0.17
Frequency of shots (shot/s) 3 3
J.K.N. Murthy, B. Venkataraman / Surface & Coatings Technology 200 (2006) 2642–2652 2643
WC with varying amounts of Co content and Cr3C2–
25(NiCr) coatings deposited by different thermal spray
processes has been studied by various researchers [7,9,11].
It has been reported that the abrasive wear rate for the
cermet coatings is controlled by several factors like the
morphology of the starting powder, the size and distribution
of the carbide particles, hardness of the carbide particles
relative to the abrasive, properties of the matrix and its
volume fraction and the coating process, which determines
the coating characteristics like the phases, density, macro-
hardness and the residual stresses [11–13]. It has been
found that the Cr3C2–25(NiCr) coating has less wear
resistance compared to WC–Co system [14–16]. Several
mechanisms of material removal have been proposed—(i)
extrusion of the binder phase and removal by plastic
deformation and fatigue, (ii) undermining of the particles
and subsequent particle pull-out, (iii) microcutting, (iv)
carbide grain fracture and (v) delamination of the coating
[12,13,17]. In the present study, the abrasive wear
behaviour of WC–10Co–4Cr and Cr3C2–20(NiCr) coatings
deposited by HVOF and pulsed DS processes have been
compared. The wear performance of these coatings is also
compared with the hard chrome coating.
2. Experimental
2.1. Material
The WC–10Co–4Cr (Metco 5847) and Cr3C2–20(NiCr)
(Diamalloy 3007) powders were coated on mild steel
substrates. The coating powder characteristics as supplied
by the manufacturer are given in Table 1. Prior to the
coating, the mild steel substrates of dimensions 70�25�10
mm thick were ultrasonically cleaned with acetone and grit
blasted using Al2O3 grits on the 70 by 25 mm coating face
and again cleaned ultrasonically with acetone and dried.
The grit blasted substrate was held suitably in a fixture and
the coating deposition was carried out with the samples in
the stationary condition and the gun traversing to and from
to obtain the desired coating thickness. The above powders
were deposited by the high velocity processes namely
HVOF and DS. The spraying conditions adopted for the
two processes are given in Table 2. The coating
Table 1
Coating powder characteristics
Characteristic Coating powder
SM 5847 D-3007
Composition 10 wt.% Co 80 wt.% Cr3C2
4 wt.% Cr 20 wt.% (Ni20Cr)
86 wt.% WC
Particle size �53+11 Am �45+5.5 AmShape Mostly spherical Irregular
Manufacturing route Agglomerated/sintered Clad
thicknesses for the two coatings were between 250 and
350 Am.
2.2. Characterization
The following characterization tests were carried out:
1. X-ray diffraction (XRD) analysis of the powders and
the coating was done using a Phillips PW 1320
diffractometer with Cu-Ka radiation operated at 40
kV and 25 mA.
2. Surface roughness measurements of the coated surface
using a surface roughness tester (Make: Mitutoyo,
Model: Surftest 211). The cut-off length was 0.8 mm.
An average of five readings is reported.
3. Scanning electron microscopy (SEM) of the sectioned
and polished surface of the coating and also the worn
surface was obtained using a Phillips make SEM.
4. Porosity measurements were done using a Leitz micro-
scope fitted with a Biovis image analyzer, on the
sectioned and polished surface of the coating. Ten
readings were obtained and the average is reported.
5. Microhardness measurements were done on the sectioned
and polished surface of the coating with a Vickers
indenter at a load of 300 g using a Leica Microhardness
Tester (Model: VMHT Auto). An average of five
readings is reported.
6. Scratch test was also performed using a pin-on-disc
tribometer. The tribometer was modified to hold a
Vickers microindenter to perform the scratch test. The
test was done on a polished surface with a surface
roughness of Rab0.15 Am. Before each test, the coated
surface and the indenter tip were cleaned with acetone.
The axis of the indenter was normal to the coating
surface and a constant load of 10 N was applied during
Table 3
Test conditions
Normal load (N) 45
Wheel (rpm) 201
Total sliding distance (m) 8657
Wheel surface speed (m s�1) 2.4
Abrasive material Silica
Particle size range (Am) 150–300
Feed rate (kg/h) 19.32
J.K.N. Murthy, B. Venkataraman / Surface & Coatings Technology 200 (2006) 2642–26522644
the scratch. The scratch length was approximately 15
mm. The scratch test was done to simulate the behaviour
of the coating material to a single asperity contact.
2.3. Abrasion wear test
The coated samples of 70�25�10 mm dimensions were
tested using the well known dry abrasive rubber wheel
tester. The coated sample was mounted firmly in the sample
holder and was allowed to press against the rim of the
rubber wheel with the desired normal force by applying a
known dead weight using a lever arrangement. The dry
silica sand was then allowed to fall freely between the
wheel and the coated surface while the rubber wheel was
rubbing against the coated surface. The abrasive particles
used were not re-cycled. Fig. 1 shows the SEM micrograph
of the abrasive particles used in the present study. Prior to
the test, the coated sample was ultrasonically cleaned with
acetone, dried and weighed using an electronic weighing
balance (Make: Sartorius) with an accuracy of 0.01 mg.
The coating mass loss was measured at every 10-min
interval. The total duration of the test was 60 min. The
mass loss obtained was normalized with the coating density
to obtain the volume wear loss. The test conditions
followed are given in Table 3. The coating density values
were estimated by the coating weight gain method in which
the mass of the coating deposited was normalized with the
coating volume. The coating density values so obtained
ranged from 11,000 to 12,500 kg m�3 for WC–10Co–4Cr
and 6100 to 7000 kg m�3 for Cr3C2–20(NiCr) coating. The
average coating density values of 12,000 and 6400 kg m�3
for the tungsten carbide-based coating and chromium
carbide-based coating were taken, respectively. For com-
parison, abrasive wear test was also done on hard chrome
plating deposited on mild steel substrate. The hard chrome
plating was done by an electroplating unit using 1%
sulphuric acid bath. Standard procedures normally followed
Fig. 1. The SEM micrograph of the silica abrasive particles.
for commercial applications were adopted. The plating
thickness was approximately 175 Am.
3. Results and discussion
3.1. X-ray diffraction analyses
The X-ray diffraction pattern for the coating powders are
shown in Fig. 2a and b, respectively. In the case of WC–
CoCr, WC and Co were detected, whereas for Cr3C2–
20(NiCr), Cr3C2 and the binder NiCr were the major phases
identified. Fig. 2c shows the XRD pattern for as-sprayed
WC–CoCr coating by HVOF process. It shows partial
decarburisation of tungsten carbide to di-tungsten carbide
(W2C). Similar partial decarburisation occurs during DS
process as well shown in Fig. 2e. Decarburisation of WC to
W2C during the deposition process has been observed [8,18].
Distinct Co peak was not present, however, broadening of
the peak was observed. Probably the binder material is
present in amorphous/nanocrystalline form, presumably due
to the high cooling rates (typically 106–107 K/s) occurring in
such deposition processes upon impact of the molten
particles on the target material/sample. Also, the occurrence
of W and C in the binder phase due to the dissolution of WC
in molten Co during the deposition process has been reported
[18]. The XRD pattern for the as-sprayed Cr3C2–20(NiCr)
coating by HVOF and DS processes are shown in Fig. 2d and
f, respectively. Both show diffused X-ray diffraction patterns
with a number of overlapping diffraction lines of carbides—
Cr3C2, Cr7C3 (formed by decarburisation of Cr3C2) and
binder NiCr as reported earlier [9,19,20] for the composition
Cr3C2–25(NiCr). The only notable difference is the amount
of retained Cr3C2 phase, which is slightly higher for the
HVOF-sprayed coating. It has been reported that decarbur-
isation of Cr3C2 to Cr7C3 or Cr23C6 does not have a
detrimental effect on the wear resistance of the coating [9].
However, in case of WC-based coatings, the decomposition
may deteriorate the wear properties of the coating due to the
formation of brittle carbides and oxy-carbides [8].
3.2. Characterisation of the coatings
3.2.1. Microstructural characterisation
The SEM micrographs of the coating powders used are
shown in Fig. 3. The WC–CoCr powder, Fig. 3a, has mostly
Fig. 2. X-ray diffraction patterns of WC–CoCr and Cr3C2–20(NiCr) (a) and (b) coating powders, (c) and (d) HVOF coating, and (e) and (f) DS coating.
Fig. 3. SEM micrographs of the coating powders. (a) WC–10Co–4Cr and (b) Cr3C2–20(NiCr).
J.K.N. Murthy, B. Venkataraman / Surface & Coatings Technology 200 (2006) 2642–2652 2645
Table 4
Coating characteristics
Characteristics Coating
WC–CoCr Cr3C2NiCr
HVOF DS HVOF DS
Surface Roughness,
Ra (Am)
3.66F0.19 4.5F0.28 2.86F0.18 4.38F0.33
Porosity (%) 2.1F1.1 1.38F0.3 1.3F0.6 0.65F0.3
Microhardness
(HV0.3)
836F30 1096F50 880F30 894F35
Macrohardness
(HV10)
524F25 1007F35 715F20 810F25
Indentation fracture
toughness
(MPa m1/2)
3.1F0.4 4.12F0.4 2.77F0.3 3.4F1
J.K.N. Murthy, B. Venkataraman / Surface & Coatings Technology 200 (2006) 2642–26522646
spherical particles, whereas, Cr3C2–20(NiCr) powder par-
ticles (Fig. 3b) have irregular shape. The surface roughness
values of the as-sprayed samples are given in Table 4. The
DS-sprayed coatings in general resulted in slightly higher
surface roughness possibly due to the slightly higher particle
velocity, which causes more particle deformation after impact
[21]. The porosity measurements for the tungsten carbide-
based and chromium carbide-based coatings deposited by the
high velocity processes are given in Table 4. It shows that the
DS coating results in a slightly lower porosity than that of the
HVOF coating. This is evident from the SEM micrograph of
Fig. 4. SEM micrographs of the transverse section of the coatings. (a) and (b) are
are for Cr3C2(NiCr) coated by HVOF and DS processes, respectively.
the transverse section of the coatings deposited by HVOF and
DS processes shown in Fig. 4. The DS-coated samples appear
to be denser compared to the HVOF-coated samples.
3.2.2. Hardness measurements
The microhardness measurements show that DS-coated
samples result in slightly higher hardness values (Table 4).
The WC–CoCr coating deposited by HVOF process resulted
in the lowest microhardness, possibly due to the higher
percentage of porosity. The macrohardness measurements on
the coating (top surface) using Vickers indenter with a normal
load of 10 kg was also carried out as it gives an indication of
the denseness of the coating. The hardness values are given in
Table 4. The macrohardness values were less than the
microhardness values. Such differences have been reported
earlier in case of thermally sprayed coatings [22]. This has
been attributed to the planar pores (pores parallel to the
coating–substrate interface) and the microcracks within the
coating. During microindentation the deformation is highly
localized, whereas in indentations at higher loads, the
influence of planar pores and cracks are more pronounced,
thus resulting in lower macrohardness values. This is
reflected in the hardness values for HVOF deposited WC–
CoCr coating which has higher porosity content that causes
substantial reduction in macrohardness value. The SEM
micrographs of the indentations produced by macroindenta-
for WC–CoCr coated by HVOF and DS processes, respectively; (c) and (d)
Fig. 5. SEM micrographs of Vickers indentations on the coated surface at a load of 98.4 N. (a) and (b) are for WC–CoCr coated by HVOF and DS processes,
respectively; (c) and (d) are for Cr3C2(NiCr) coated by HVOF and DS processes, respectively.
J.K.N. Murthy, B. Venkataraman / Surface & Coatings Technology 200 (2006) 2642–2652 2647
tion are shown in Fig. 5. It can be seen that edge cracks are
formed in all the coatings. Such cracks are normally observed
in thermally sprayed coatings as the cracks at the edges are
formed by the coalescence of microcracks originating at the
pores. In the case of HVOF WC–CoCr (Fig. 5a) coating, the
edge crack intensity surrounding the indentation was higher
probably due to the higher porosity content in the coating.
3.2.3. Indentation fracture toughness measurements
The indentation technique was used to obtain the fracture
toughness of the coatings using a Vickers indenter. The
Fig. 6. The indentation cracks induced in WC–CoCr
indentation was carried out on the transverse section of the
coating in the mid-plane region to minimize the edge and
interface effects [23]. The indenter was loaded such that one
of the horizontal diagonals was parallel to the coating–
substrate interface. A load of 2 kg was applied for a dwell
time of 25 s at a rate of 25 Am/s. Figs. 6 and 7 show the typical
indentations on the transverse section with in-plane cracks for
WC–CoCr and Cr3C2–20(NiCr) coatings respectively. In the
thermally sprayed coatings, the cracks parallel to the coating–
substrate interface are more easily formed in comparison to
the perpendicular direction [24,25]. This has been attributed
coating: (a) HVOF coating and (b) DS coating.
Fig. 7. The indentation cracks induced in Cr3C2–20(NiCr) coating: (a) HVOF coating and (b) DS coating.
Fig. 8. The indentation crack propagation path in (a) DS-coated WC–CoCr
and (b) HVOF-coated Cr3C2–20(NiCr).
J.K.N. Murthy, B. Venkataraman / Surface & Coatings Technology 200 (2006) 2642–26522648
to the characteristics of the thermally sprayed coatings [24].
The coatings are built to the desired thickness by the
deposition of molten or semi-molten particles in the form of
plate-like structure called splats. The weak bonding between
the splats results in such an anisotropic crack formation. In
some indentation tests, the cracks were not formed. As can be
seen (Figs. 6 and 7) in most of the cases, the cracks were not
initiated at the corners of the indentation where the stress is
highest but are seen originating from the sides of the
indentation, this has been attributed to the non-uniform
microstructure in the as-sprayed coatings [23]. The crack
length from the center of the indent, c, was used for
determining the fracture toughness of the coatings. The
fracture toughness values of the coating were calculated
according to the method suggested by Evans and Wilshaw
[26]. The length of the cracks was measured from the SEM
images. The Table 4 shows fracture toughness values
(average of 10 readings) obtained for the coatings. The
fracture toughness values showed some scatter for the two
coatings indicating non-homogeneous coating microstruc-
tural features. Such a variation in the fracture toughness
values in thermally sprayed coatings has been observed [23].
In case of WC–CoCr HVOF coating, the crack lengths could
not be measured accurately due to the interference of the
porosity. WC–CoCr coating had marginally higher fracture
toughness compared to Cr3C2–20(NiCr) coatings. Further,
examination of the crack features show that the indentation
crack propagates along a region between the carbide particles
and the binder phase in case of both WC–CoCr and Cr3C2–
20(NiCr) coatings as shown in Fig. 8a and b, respectively.
Similar observation on the indentation crack propagation has
been reported earlier in case of WC–CoCr coating [25].
Further investigation is currently being carried out in our
studies to understand the metallurgical characteristics influ-
encing the crack propagation path.
3.2.4. Scratch behaviour
To understand the wear behaviour when a single abrasive
particle rubs against the coating surface, a pin-on-disc
tribometer was used. It was modified to hold a Vickers
microindenter to perform the scratch test at a constant load.
This may be an over-simplification considering the fact that
the indenter used is diamond whose hardness is significantly
higher compared to the coating materials and acts as a hard
abrasive resulting in severe wear conditions. However, this
test may give a qualitative understanding of the response of
the coating materials under identical test conditions
employed. Fig. 9 shows the groove morphology formed
Fig. 9. SEM micrographs of the surface scratched by Vickers microindenter at 10 N load. (a) and (b) are for WC–CoCr coated by HVOF and DS processes,
respectively; (c) and (d) are for Cr3C2(NiCr) coated by HVOF and DS processes, respectively.
Fig. 10. Volume wear loss as a function of time. The filled symbols
represent HVOF process and the unfilled symbols DS process.
J.K.N. Murthy, B. Venkataraman / Surface & Coatings Technology 200 (2006) 2642–2652 2649
along with the wear debris scattered on either side during the
scratch test. There is a good correlation between the scratch
width and the hardness of the coating. As the indenter
scratches the surface, material is displaced to the sides and
detached. Some of the displaced material will be accom-
modated below the indenter by the porous volume and also
possibly by the surrounding elastic strain field. The displaced
volume of material by the indenter due to ploughing forms
side ridges, which may consist of the binder phase with few
carbide particles. Due to the near zero hydrostatic compres-
sive forces on the side of the indenter at the free surface,
plastic strain in the material flowing on the sides reaches a
high value [22,27]. Beyond a certain critical strain, the
material detaches to form debris. Here, the material detach-
ment is observed to occur in a single pass for both the coated
samples tested. The morphology of the debris also shows the
evidence of microcutting when the indenter (hard abrasive)
scratches the surface. The displacement of the material to the
sides is slightly more in the case of DS-coated samples
possibly due to the slightly higher density of the DS coatings.
3.3. Abrasive wear behaviour
The incremental volume wear loss of the coating as a
function of time and the cumulative wear loss of the coatings
are shown in Figs. 10 and 11, respectively. It shows that WC–
CoCr coating has a lower volume wear loss compared to
Cr3C2(NiCr) coating. The results show that the abrasive wear
volume was slightly lower in the case of DS coating
compared to that of HVOF coating. The SEM micrographs
of the worn surface of the coatings are shown in Fig. 12. In the
case of WC–CoCr coating (Fig. 12a and b), mechanism of
wear was by selective removal of the binder caused probably
by plastic deformation and fatigue due to the repeated action
Fig. 11. The cumulative volume wear loss of the thermally sprayed coatings in comparison with the hard chrome plating.
J.K.N. Murthy, B. Venkataraman / Surface & Coatings Technology 200 (2006) 2642–26522650
of the abrasive particles followed by the undermining of the
carbide particles resulting in their eventual pullout. Some
evidence of microcutting may also be noticed indicating the
removal of the binder phase by this mechanism. Also, very
Fig. 12. SEM micrographs of the worn surfaces. (a) and (b) are WC–CoCr coating
coating by HVOF and DS processes, respectively. The direction of abrasion is fr
little carbide grain fracture was observed. In the case of
Cr3C2(NiCr) coating (Fig. 12c and d) in addition to the above
mechanism of wear, material removal by delamination in
both HVOF- and DS-coated samples are observed. The
by HVOF and DS processes, respectively; (c) and (d) are Cr3C2–20(NiCr)
om top to bottom.
J.K.N. Murthy, B. Venkataraman / Surface & Coatings Technology 200 (2006) 2642–2652 2651
transverse section of the worn region was examined to
observe the sub-surface damage in both the coating systems.
The transverse section of the central area of the worn region
along the direction of sliding is shown in Figs. 13 and 14. In
the case ofWC–CoCr (Fig. 13), there is little evidence of sub-
surface damage. However, in very few regions, some
microcracks were observed below the surface running
parallel to the surface. With Cr3C2–NiCr coating (Fig. 14),
material removal by sub-surface crack formation and its
propagation to the free surface to form large chunks of debris
is evident. Such sub-surface cracking was more extensive for
Cr3C2–20(NiCr) coating. Such wear behaviour was common
to both the high velocity coatings as seen in Fig. 14.
The higher wear resistance of WC–CoCr coating over
Cr3C2–20(NiCr) may probably be due to the higher fracture
strength and better adhesive strength of the CoCr matrix
with the carbides compared to the NiCr matrix. The slightly
better indentation fracture toughness of the WC–COCr
coating observed may support this behaviour. The delami-
nating-type cracking has been observed earlier in case of
Cr3C2–25(NiCr) coating, indicating that the matrix phase is
not acting as an efficient toughening phase [14,28]. Such
Fig. 13. SEM micrographs showing SE images of the transverse section of
the worn region for WC–CoCr coating by (a) HVOF process and (b) DS
process.
Fig. 14. SEM micrographs showing SE images of the transverse section of
the worn region for Cr2C3–20(NiCr) coating by (a) HVOF process and (b)
DS process.
type of wear results in the removal of much higher material
thus leading to higher wear rates. This may explain the
better wear resistance offered by HVOF WC–CoCr coating
even in the presence of slightly higher porosity. Also, the
higher hardness of WC grains compared to the Cr3C2 grains
will enable the former to withstand higher loads which
results in fracture of less number of WC grains.
The residual stress measurements by X-ray diffraction
technique was also carried on the WC–CoCr coating. The
Bragg reflection from the (256) planes of WC was used for
determining the residual stresses. More detailed description
on the residual stress measurements is given in our earlier
paper [29]. The results show that compressive residual
stresses are induced in the coating deposited by both the
processes. In fact, higher compressive residual stresses were
induced in the DS coating (�104F20 MPa) compared to the
HVOF coating (�25F10 MPa) probably due to the higher
particle velocities in the former process. This may also
explain the lower wear rate observed in the DS coating
where the compressive stresses impedes or delays the crack
initiation and propagation, thus generating less wear debris.
J.K.N. Murthy, B. Venkataraman / Surface & Coatings Technology 200 (2006) 2642–26522652
Abrasive wear test of hard chrome plating was also carried
out for comparison with the thermally sprayed coatings. The
hardness of the hard chrome plating was 871HV0.3 (an
average of five readings). The results (Fig. 11) show that the
carbide-based thermally sprayed coatings have better abra-
sive wear resistance compared to the hard chrome plating. It
is also important to mention that the wear test on the hard
chrome plating had to be discontinued after 20 min as the
plating got completely peeled off the substrate at the test
region under identical test conditions, indicating the poor
adhesion of the plating to the substrate. The tensile stresses
induced during the plating process and the microcracks
present may also lead to its rapid wear suggesting the superior
performance of both the carbide-based coatings against
abrasion.
4. Conclusions
1. WC–10Co–4Cr has better abrasive wear resistance
compared to Cr3C2–20(NiCr) coating possibly due to
the higher hardness of WC particles and better matrix
properties of the CoCr binder material. The indentation
fracture toughness of WC–CoCr coating was slightly
better than that for Cr3C2–20(NiCr), which may also
cause the wear resistance to improve in the former.
2. The coatings deposited by the DS process had slightly
improved wear resistance compared to the HVOF process
possibly due to the higher density and compressive
residual stresses induced in the DS coating.
3. Both the thermally sprayed coatings had superior wear
performance in comparison to the hard chrome coating.
Acknowledgements
The authors wish to thank the Director, DMRL, for
granting permission to publish the paper. Also, the support
provided by the Electron Microscopy group and the X-ray
group of DMRL is acknowledged.
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